CA1136905A - Superalloy composition and process - Google Patents
Superalloy composition and processInfo
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- CA1136905A CA1136905A CA000324088A CA324088A CA1136905A CA 1136905 A CA1136905 A CA 1136905A CA 000324088 A CA000324088 A CA 000324088A CA 324088 A CA324088 A CA 324088A CA 1136905 A CA1136905 A CA 1136905A
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Abstract
SUPERALLOY COMPOSITION AND PROCESS
ABSTRACT OF THE DISCLOSURE
A nickel superalloy composition and process for producing novel, high strength articles of the alloy are described. The composition is based on the Ni-Al-Mo system. Outstanding properties are obtained by producing the alloy in powder form, compacting the powder and causing directional secondary recrystallization in the compacted powder. The resultant articles have high temperature properties which are superior to those of known commercial superalloys.
ABSTRACT OF THE DISCLOSURE
A nickel superalloy composition and process for producing novel, high strength articles of the alloy are described. The composition is based on the Ni-Al-Mo system. Outstanding properties are obtained by producing the alloy in powder form, compacting the powder and causing directional secondary recrystallization in the compacted powder. The resultant articles have high temperature properties which are superior to those of known commercial superalloys.
Description
1~3690S
BACK~ROUND OF THE INVENTION
The invention described herein was made in the course of a contract with the Defense Advanced Research Projects Agency of the U.S. Government.
Field of the Invention - This invention relates to the field of the nickel base superalloys and processes to enhance the properties of nickel base superalloys.
Descr~iption of the Prior Art - U.S. Patent,2,542,962 to Kinsey discloses a Ni-Al-Mo alloy with broad ranges which encompass the alloy of the present invention. Kinsey does not discuss the possibility of heat treatments~
U.S. Patent 3,617,397 to Maxwell, assigned to the present assignee, discloses a cast nickel base superalloy containing 8% Al and 18% Mo, however, the patent does not disclose the specific processing steps which form a part of the present invention.
U.S. Patent 4,012,241 and U.S. Patent 4,111,723 which are both assigned to the present assignee, relate to directionally solidified eutectic compositions in the Ni-Al-Mo system.
U.S. Patent 3,975,219 to Allen et al, also assigned to the present assignee, describes a directional recrystalli-zation process in which a hot worked nickel base superalloy article is recrystallized in a thermal gradient to produce an elongated grain microstructure.
~13~5~05 SUMMARY OF THE I NVE~TI ON
The present invention concerns an alloy composition and process which can be employed together to produce an article having exceptional high temperature mechanical properties. The broad composition range is 5-13% Al, 5-30/O Mo, balance nickel. Preferably, additions of Ta are made. The alloy composition is prepared in the form of a homogeneous powder which is compacted to form a fine grain article whose grain size is stabilized by gamma prime parti-cles at the grain boundaries. This fine grained article is passed through a thermal gradient so as to dissolve the gamma prime particles and to permit grain growth to occur through a directional secondary recrystallization effect. The grains which result in the article are elongated in the direction of relative motion between the article and the thermal gradient.
- The article is then solution treated and aged to achieve mechanical properties at elevated temperatures which are greatly in excess of t~e properties of currently used nickel base superalloys.
More specifically, according to the invention, there is provided a method for producing a high strength aligned grain superalloy article from an alloy which contains from 5-10% Al and from 8-21% Mo, balance nickel, including the steps of: a~ forming the alloy into a powder, b. compacting the powder alloy into an article; c. heating the article at a temperature between 2200F for a period of time in excess of one hour, and the gamma prime solvus so as to form gamma prime particles at the grain boundaries; d. progressively and selectively heating the article by causing relative motion between the article and a thermal gradient, said thermal gradient having a hot end temperature which lies between the gamma prime solvus temperature and the incipient melting ~36~05 temperature, so that the gamma prime particles at the grain boundaries dissolve and directional grain growth occurs in the direction of relative motion between the article and the thermal gradient; whereby said superalloy article has an aligned grain structure and at a temperature of 1800F a stress in excess of 33 ksi is required to produce more than 1% of elongation in a time period of 100 hours.
The wrought recryst,allized powder metallurgy article according to the invention comprises elongated aligned grains of the gamma phase which contain a fine dis-persion of the gamma prime phase, said article having a composition of 5-10% Al, 8-21% Mo, up to 12% Ta, up to 1% Y, balance essentially Ni; said article being capable of with-standing stresses of up to 33 ksi at 1800F for 100 hours without undergoing more than 1% elongation.
The foregoing and other objects, features and ad-vantages of the present invention will become more apparent in the light of the following detailed description of prefer~
red embodiments thereof as illustrated in the accompanying drawings.
- 3a -BRIEF DESCRIPTION OF THE DRAWINGS
Fig. 1 is a photomicrograph, after extrusion, of the invention alloy;
Fig. 2 is a photomicrograph of the invention alloy after heat treatment in preparation for directional secondary recrystallization;
Fig. 3 is a photomicrograph of the invention alloy after directional secondary recrystallization;
Fig. 4 is a photomicrograph after solution heat treatment and air cooling;
Fig. 5 is a creep curve showing time to 1% elonga-tion for various combinations of stress and temperature;
Fig. 6 is a creep curve showing elongation as a function of time at a specific temperature and stress;
Fig. 7 is a photomicrograph of the invention alloy containing a low Ta level after solution treatment;
Fig. 8 is a photomicrograph of the invention alloy (lcw Ta) after solution treatment and aging; and Fig. 9 is a photomicrograph of the invention alloy (high Ta) after solution treatment and aging.
DESCRIPTION OF PREFERRED EMBODIMENTS
Unless otherwise noted~ all percentages listed herein are weight percentages.
The present invention relates to a superalloy article of specific composîtion and microstructure which is the product of a specific process. The article has exceptional mechanical properties at elevat~d temperature.
The basic alloy is a simple one based on the nickel-aluminum-molybdenum system. The broad composition ranges are from 5-13% Al, from 5-30% Mo, balance nickel. The preferred ranges are 5-10% Al and 8-21% Mo. These ranges overlap certain other alloys, in particular the alloys disclosed in U. S. Patent 3,655,462. The present inven-tion, however, achieves une~pected mechanical properties through the combination of the alloy and the processing sequence.
The basic ternary alloy has good short-term mechan-ical properties at elevated temperatures but suffers from microstructural instabilities which cause a marked deterioration in properties after long-term exposures at elevated temperatures. These instabilities include the formation of massive gamma prime particles at the grain boundaries and the precipitation of Mo from the gamma phase. However, for certain applications the basic alloy may be adequate. The addition of Ta to the alloy in amounts of up to about 12% tends to stabilize the microstructure and improve the long-term mechanical properties at elevated temperatures. At least about 4% Ta appears necessary to significantly stabilize the microstructure. Ta is a relatively dense element and its incorporation into the alloy raises the density, consequently a preferred Ta range is from about 4 to 113~i~05 about 10%. The Ta is observed to largely substitute for Al and when it is added the Al content may be reduced in proportion to the atomic percent oE Ta added.
The major phases present in alloys of these compo-sitions are the gamma and gamma prime phases. The gamma phase (nickel solid solution~ is the matrix phase in which are found discrete particles of the gamma prime phase (Ni3Al) and other phases such as Ni2Mo which are present as extremely fine dispersions. The gamma prime phase is present in amounts of 40 to 70 volume percent.
The alloys have a gamma prime solvus temperature, that temperature above which the gamma prime phase dissolves into the gamma phase, which ranges from about 2320-2440F. The solidus and liquidus temperatures of the alloys are very close together and range from about 2350-2470F. The solidus temperature preferably exceed the gamma prime solvus temperature by at least 30F so th~t the alloy may be solution treated without incipient melting.
The microstructure achieved by the process to be described below is polycrystalline microstructure and it is well known that grain boundaries often adversely affect the high temperature properties of materials. Such adverse eEfects can be minimized through the addition o~
certain interstitial elements such as C and B. These interstitial elements are believed to segregate to the grain boundaries. Other noninterstitial elements such li3690S
as Y, La and Ce produce similar effects. Particularly outstandin~ results have been obtained with the addition of about 150 ppm of Y and to a lesser extent with the addition of abou-t .05% C to the present alloys. The other elements have not proven to be as beneficial, although they have not been extensively investigated.
Y, in amounts of up to 1%, and C in amounts of up to about .1% are the preferred additions. B may be added in amounts of up to about .05% and Zr in amounts of about up to .1%. Other elements such as Ce and La may ^
be substituted in part for the Y addition.
Part or all of the Ta may be replaced by an equiatomic amount of Cb, Ti or W. Likewise, part of the ; Mo may be replaced by an equiatomic amount of W or Re.
For certain applications, up to 5% Cr and up to 5% Co might be added.
The preferred alloy consists of from 5-10% Al, from 8-21% Mo, from 4-12% Ta, up to 1% Y and up to .1% C.
The alloy previously described must be prepared in the form of powder. This powder preparation process is preferably one which involves the solidification of molten metal and most preferably involves a high cooling rate solidification process. Such a high cooling rate helps to insure compositional homogeneity of the resulting powder. All the work in connection with this invention has employed a centrifugal atomization apparatus such as that shown in U. S. Patents 4,025,249; 4,053,264 and 1~36905 4,078,873. The centrifugal atomization process involves pouring the molten metal to be atomized on a rapidly rotating disk. The molten metal is thrown off the edge of the disk in a finely divided form. The particular process employed used an extremely high rate of disk rotation, up to 35,000 rpm and also employed auxiliary gas cooling which was provided by a flowing curtain o~ helium gas surrounding the rotating disk through which the atomized metal passed. The powder produced by this process had an average particle size of 70 microns and the rate of cooling was about 105-10 C per second. All of the experimental work described herein employed powder of the type described. Coarser powders or more slowly solidified powders may require longer times at elevated temperatures to achieve compositional homogeneity. The powder is produced and maintained in a low oxygen environment so that the resultant powder contains less than about 50 ppm of oxygen. The nitrogen level is also maintained at or below 50 ppm.
The powder was compacted under conditions which caused interparticle bounding and produced a void-free article of sub-stantially theoretical density. Experimental work employed hotextrusion as the compaction technique. Stainless steel con-tainers were evacuated, filled with the powder to be compacted and sealed. The filled containers were preheated and extruded.
Extrusion was conducted at temperatures between 2200 and 2300F, and extrusion ratios of from 6:1 to 43:1 have been successfully employed. Fig. 1 shows the typical microstructure of the ~3690:~
extruded powder. The exact nature of the compacting process does not appear to be critical to the invention. other compaction techniques including those performed at low temperatures such as explosive compaction and those performed at elevated temperatures where there is no significant metal flow such as hot isostatic compaction may be employed. It is important, however, that during the compaction process the metal temperature is maintained below the gamma prime solvus, so that grain growth is avoided and a fine grain microstructure is maintained. The compacted grain size is less than the particle size of the starting material. The gamma prime second phase is effective to eliminate grain boundary migration and thus maintain the initial grain size.
The material at this point has a fine uniform grain size and it is observed that the microstructure is duplex.
The gamma phase grains have a grain size less than the start-ing particle size. Finer particles of the gamma prime phase are located at the grain boundaries and extremely fine particles of gamma prime are located within the gamma phase matrix. It is desirable that the amount of gamma prime at the grain boundaries be maximized and this can be achieved by a heat treatment step at a ~ .
113~i90S
temperature bet~een about 2200F and the gamma prime solvus temperature for a period of time in excess of about one hour. The treatment used in the ~ork described herein was 4 hours at 2300F. The heat treatment will greatly increase the fraction of the gamma prime phase found at the grain boundaries which is important in achieving controllable and reproducible grain growth in subse~uent process steps. This heat treatment step forms a part of the preferred embodiment.
Of course, this heat treatment step may be combined with other process steps, for example, if hot extrusion is used to compact the powder, the extruded material may be slowly cooled so that the period of time spent between the gamma prime solvus temperature and about 2200F is sufficient to permit gamma prime phase growth at grain boundaries.
The material at this point in the process sequence has a uniform fine grain size and consists of grains of the gamma phase which contain very fine gamma prime particles on the order of .5-3 microns and larger gamma prime phase particles on the order of 5-20 microns are located at the grain boundaries. This microstructure is shown in Fig. 2. The light colored phase is the gamma prime phase. The fine grain size of ~his material is stable at temperatures up to about the gamma prime solvus temperature. More importantly, this fine grain material displays superplastic behavior and may be ~3~i905 deformed at temperatures from about 1800F to the gamma prime solvus temperature with significant ductility and a low flow stress.
To illustrate this behavior, a one-inch bar of extruded RSR 104 material (composition shown in Table I) was heat treated for 4 hours at 2300F and then hot rolled. Several samples were hot rolled at temperatures oE 2200 and 2300F. The samples were red~ced from a starting thickness of 1.0 inch to a final thickness of .030 inch. Between 12 to 15 passes were employed to achieve the total reduction with a reduction per pass of from .030 to .080 inch. After each pass, the material was returned to a furnace maintained at 2200-2300F for a period of time sufficient to bring the material tempera-ture back up to the starting temperature. No difficulty was encountered in hot rolling this material and this success is somewhat surprising in view of the notorious difficulties encountered in hot working conventional superalloys. This behavior has important commercial ~0 implications since it will make possible fabrication of intricate shapes at low cost. The deformation step is of course optional and not required.
The next step is a directional grain growth step which makes use of the phenomena known as secondary recrystallization. An apparatus is required which can produce a large temperature difference over a small distance (a steep gradient). The hot end of the gradient `` ` 1~3~905 exceeds t~e ga~a pri~.e solvus t~,perature DUt iS less th2n the incipient melting te~perature. Th~ article is sLo~71y ~oved relative to the thermal gr2dient so that it is progressively and selectively heated to ~ te~.perature bet~,een the ga~.a primQ sol~Jus temperatùre and the incipient r,elting temperature. As the article is heated within this temperature range, the ga~.a pri~e particl~s which have stabilized the grain boundarie~ dissolve permitting grain gro~th to occur. Grain grot~th occurs, the driving force being a reduction in grain bou~dary area. The grains will grow in the direction of relat.ive motion between the article and the ther~al gradient and the resultant grains will be elongated having a major dimJension ~7hich is at least ten times the minor dimension.
The thermal gradient should be as steep as possiDle in order to ~aximiz2 the directiona,lit~- o~ the grains and mini~i~e the possibi'ity of laterzl grain gro~.7th and resultant transverse grain boundaries. The ther~al gradient should be at least 20F per inc'n ~.ea~ured at the ga~lm2 prime solvus, but prererably gr~ater. The experimenta~ work reported herein was done with a thermal 0~
/; ,,,.-j,fgradient of 150 to 300F per inch1. The Oradient T~as obtained by using induction heatinD with a graphit2 susceptor through ~hich the article passes The raLe at which the article c2n be ~oved rel2tive -o the ther~al gradient 2ppears to be on the order o~ 2bcut one-hal inch per hour. For obvious co~.~.erci~l re~sons, it is 3 ~
~36905 desirable to move the workpiece relative to the gradient at as rapid a rate as possible.
Following the directional secondary recrystallization step, the material will have a rather coarse gamma prime microstructure, within the elongated grains, as shown in Fig. 3. This is a consequence of the slow cooling rate through the gamma prime solvus temperature which permits a large amount of gamma prime phase nucleation and growth. It is preferred that as a final step the entire article be solution heat treated at a temperature between the gamma prime solvus temperature and the incipient melting temperature for a period of time suffi-cient to produce a solid solution, then rapidly cooled and aged to produce a refined gamma prime structure.
The conditions experimentally used were 4 hours above the gamma prime solvus followed by an air cool to room temperature and the article was then reheated to 1950F
for 4 hours and 1600F for 12 hours. Those skilled in the art will appreciate that these conditions may be ~0 widely varied and that they may be combined with other operations such as coating operations which involve heating the article. The effect of the heat treatment is to produce a very refined dispersion of gamma prime part;cles. Fig. 4 shows the microstructure after the aging step. The refined microstructure contributes to the improved mechanical properties.
Table I gives the compositions of three experimental 1~3~905 alloys, processed according to the invention, denoted as RSR alloys and two commercial prior art superalloys, PWA 1422 and PWA 1419. RSR 104 alloy is the basic ternary Ni-Al-Mo. RSR 166 is the same alloy composition as RSR 104 with the addition of 150 ppm Y. RSR 143 is a Ni-Al-Mo alloy with 6% Ta added. PWA 1422 is one of the strongest commercially used nickel base superalloys.
PWA 1419 is a prior art nickel base superalloy containing Ni, Mo, Al and C. This alloy is claimed in U. S. Patent 3,655,462. Both PWA 1422 and 1419 are used in the directional solidified form as described in U. S. Patent 3,260,505. Directional solidification is a casting pro-cess which produces a microstructure containing elongated grains. The properties discussed below were obtained from directionally solidified samples. Fig. 5 is a graph which shows the creep properties of the various superalloys in Table I. Fig. 5 shows the time required to reach 1% elongation in creep as a function of stress and temperature. Time and temperature are combined in
BACK~ROUND OF THE INVENTION
The invention described herein was made in the course of a contract with the Defense Advanced Research Projects Agency of the U.S. Government.
Field of the Invention - This invention relates to the field of the nickel base superalloys and processes to enhance the properties of nickel base superalloys.
Descr~iption of the Prior Art - U.S. Patent,2,542,962 to Kinsey discloses a Ni-Al-Mo alloy with broad ranges which encompass the alloy of the present invention. Kinsey does not discuss the possibility of heat treatments~
U.S. Patent 3,617,397 to Maxwell, assigned to the present assignee, discloses a cast nickel base superalloy containing 8% Al and 18% Mo, however, the patent does not disclose the specific processing steps which form a part of the present invention.
U.S. Patent 4,012,241 and U.S. Patent 4,111,723 which are both assigned to the present assignee, relate to directionally solidified eutectic compositions in the Ni-Al-Mo system.
U.S. Patent 3,975,219 to Allen et al, also assigned to the present assignee, describes a directional recrystalli-zation process in which a hot worked nickel base superalloy article is recrystallized in a thermal gradient to produce an elongated grain microstructure.
~13~5~05 SUMMARY OF THE I NVE~TI ON
The present invention concerns an alloy composition and process which can be employed together to produce an article having exceptional high temperature mechanical properties. The broad composition range is 5-13% Al, 5-30/O Mo, balance nickel. Preferably, additions of Ta are made. The alloy composition is prepared in the form of a homogeneous powder which is compacted to form a fine grain article whose grain size is stabilized by gamma prime parti-cles at the grain boundaries. This fine grained article is passed through a thermal gradient so as to dissolve the gamma prime particles and to permit grain growth to occur through a directional secondary recrystallization effect. The grains which result in the article are elongated in the direction of relative motion between the article and the thermal gradient.
- The article is then solution treated and aged to achieve mechanical properties at elevated temperatures which are greatly in excess of t~e properties of currently used nickel base superalloys.
More specifically, according to the invention, there is provided a method for producing a high strength aligned grain superalloy article from an alloy which contains from 5-10% Al and from 8-21% Mo, balance nickel, including the steps of: a~ forming the alloy into a powder, b. compacting the powder alloy into an article; c. heating the article at a temperature between 2200F for a period of time in excess of one hour, and the gamma prime solvus so as to form gamma prime particles at the grain boundaries; d. progressively and selectively heating the article by causing relative motion between the article and a thermal gradient, said thermal gradient having a hot end temperature which lies between the gamma prime solvus temperature and the incipient melting ~36~05 temperature, so that the gamma prime particles at the grain boundaries dissolve and directional grain growth occurs in the direction of relative motion between the article and the thermal gradient; whereby said superalloy article has an aligned grain structure and at a temperature of 1800F a stress in excess of 33 ksi is required to produce more than 1% of elongation in a time period of 100 hours.
The wrought recryst,allized powder metallurgy article according to the invention comprises elongated aligned grains of the gamma phase which contain a fine dis-persion of the gamma prime phase, said article having a composition of 5-10% Al, 8-21% Mo, up to 12% Ta, up to 1% Y, balance essentially Ni; said article being capable of with-standing stresses of up to 33 ksi at 1800F for 100 hours without undergoing more than 1% elongation.
The foregoing and other objects, features and ad-vantages of the present invention will become more apparent in the light of the following detailed description of prefer~
red embodiments thereof as illustrated in the accompanying drawings.
- 3a -BRIEF DESCRIPTION OF THE DRAWINGS
Fig. 1 is a photomicrograph, after extrusion, of the invention alloy;
Fig. 2 is a photomicrograph of the invention alloy after heat treatment in preparation for directional secondary recrystallization;
Fig. 3 is a photomicrograph of the invention alloy after directional secondary recrystallization;
Fig. 4 is a photomicrograph after solution heat treatment and air cooling;
Fig. 5 is a creep curve showing time to 1% elonga-tion for various combinations of stress and temperature;
Fig. 6 is a creep curve showing elongation as a function of time at a specific temperature and stress;
Fig. 7 is a photomicrograph of the invention alloy containing a low Ta level after solution treatment;
Fig. 8 is a photomicrograph of the invention alloy (lcw Ta) after solution treatment and aging; and Fig. 9 is a photomicrograph of the invention alloy (high Ta) after solution treatment and aging.
DESCRIPTION OF PREFERRED EMBODIMENTS
Unless otherwise noted~ all percentages listed herein are weight percentages.
The present invention relates to a superalloy article of specific composîtion and microstructure which is the product of a specific process. The article has exceptional mechanical properties at elevat~d temperature.
The basic alloy is a simple one based on the nickel-aluminum-molybdenum system. The broad composition ranges are from 5-13% Al, from 5-30% Mo, balance nickel. The preferred ranges are 5-10% Al and 8-21% Mo. These ranges overlap certain other alloys, in particular the alloys disclosed in U. S. Patent 3,655,462. The present inven-tion, however, achieves une~pected mechanical properties through the combination of the alloy and the processing sequence.
The basic ternary alloy has good short-term mechan-ical properties at elevated temperatures but suffers from microstructural instabilities which cause a marked deterioration in properties after long-term exposures at elevated temperatures. These instabilities include the formation of massive gamma prime particles at the grain boundaries and the precipitation of Mo from the gamma phase. However, for certain applications the basic alloy may be adequate. The addition of Ta to the alloy in amounts of up to about 12% tends to stabilize the microstructure and improve the long-term mechanical properties at elevated temperatures. At least about 4% Ta appears necessary to significantly stabilize the microstructure. Ta is a relatively dense element and its incorporation into the alloy raises the density, consequently a preferred Ta range is from about 4 to 113~i~05 about 10%. The Ta is observed to largely substitute for Al and when it is added the Al content may be reduced in proportion to the atomic percent oE Ta added.
The major phases present in alloys of these compo-sitions are the gamma and gamma prime phases. The gamma phase (nickel solid solution~ is the matrix phase in which are found discrete particles of the gamma prime phase (Ni3Al) and other phases such as Ni2Mo which are present as extremely fine dispersions. The gamma prime phase is present in amounts of 40 to 70 volume percent.
The alloys have a gamma prime solvus temperature, that temperature above which the gamma prime phase dissolves into the gamma phase, which ranges from about 2320-2440F. The solidus and liquidus temperatures of the alloys are very close together and range from about 2350-2470F. The solidus temperature preferably exceed the gamma prime solvus temperature by at least 30F so th~t the alloy may be solution treated without incipient melting.
The microstructure achieved by the process to be described below is polycrystalline microstructure and it is well known that grain boundaries often adversely affect the high temperature properties of materials. Such adverse eEfects can be minimized through the addition o~
certain interstitial elements such as C and B. These interstitial elements are believed to segregate to the grain boundaries. Other noninterstitial elements such li3690S
as Y, La and Ce produce similar effects. Particularly outstandin~ results have been obtained with the addition of about 150 ppm of Y and to a lesser extent with the addition of abou-t .05% C to the present alloys. The other elements have not proven to be as beneficial, although they have not been extensively investigated.
Y, in amounts of up to 1%, and C in amounts of up to about .1% are the preferred additions. B may be added in amounts of up to about .05% and Zr in amounts of about up to .1%. Other elements such as Ce and La may ^
be substituted in part for the Y addition.
Part or all of the Ta may be replaced by an equiatomic amount of Cb, Ti or W. Likewise, part of the ; Mo may be replaced by an equiatomic amount of W or Re.
For certain applications, up to 5% Cr and up to 5% Co might be added.
The preferred alloy consists of from 5-10% Al, from 8-21% Mo, from 4-12% Ta, up to 1% Y and up to .1% C.
The alloy previously described must be prepared in the form of powder. This powder preparation process is preferably one which involves the solidification of molten metal and most preferably involves a high cooling rate solidification process. Such a high cooling rate helps to insure compositional homogeneity of the resulting powder. All the work in connection with this invention has employed a centrifugal atomization apparatus such as that shown in U. S. Patents 4,025,249; 4,053,264 and 1~36905 4,078,873. The centrifugal atomization process involves pouring the molten metal to be atomized on a rapidly rotating disk. The molten metal is thrown off the edge of the disk in a finely divided form. The particular process employed used an extremely high rate of disk rotation, up to 35,000 rpm and also employed auxiliary gas cooling which was provided by a flowing curtain o~ helium gas surrounding the rotating disk through which the atomized metal passed. The powder produced by this process had an average particle size of 70 microns and the rate of cooling was about 105-10 C per second. All of the experimental work described herein employed powder of the type described. Coarser powders or more slowly solidified powders may require longer times at elevated temperatures to achieve compositional homogeneity. The powder is produced and maintained in a low oxygen environment so that the resultant powder contains less than about 50 ppm of oxygen. The nitrogen level is also maintained at or below 50 ppm.
The powder was compacted under conditions which caused interparticle bounding and produced a void-free article of sub-stantially theoretical density. Experimental work employed hotextrusion as the compaction technique. Stainless steel con-tainers were evacuated, filled with the powder to be compacted and sealed. The filled containers were preheated and extruded.
Extrusion was conducted at temperatures between 2200 and 2300F, and extrusion ratios of from 6:1 to 43:1 have been successfully employed. Fig. 1 shows the typical microstructure of the ~3690:~
extruded powder. The exact nature of the compacting process does not appear to be critical to the invention. other compaction techniques including those performed at low temperatures such as explosive compaction and those performed at elevated temperatures where there is no significant metal flow such as hot isostatic compaction may be employed. It is important, however, that during the compaction process the metal temperature is maintained below the gamma prime solvus, so that grain growth is avoided and a fine grain microstructure is maintained. The compacted grain size is less than the particle size of the starting material. The gamma prime second phase is effective to eliminate grain boundary migration and thus maintain the initial grain size.
The material at this point has a fine uniform grain size and it is observed that the microstructure is duplex.
The gamma phase grains have a grain size less than the start-ing particle size. Finer particles of the gamma prime phase are located at the grain boundaries and extremely fine particles of gamma prime are located within the gamma phase matrix. It is desirable that the amount of gamma prime at the grain boundaries be maximized and this can be achieved by a heat treatment step at a ~ .
113~i90S
temperature bet~een about 2200F and the gamma prime solvus temperature for a period of time in excess of about one hour. The treatment used in the ~ork described herein was 4 hours at 2300F. The heat treatment will greatly increase the fraction of the gamma prime phase found at the grain boundaries which is important in achieving controllable and reproducible grain growth in subse~uent process steps. This heat treatment step forms a part of the preferred embodiment.
Of course, this heat treatment step may be combined with other process steps, for example, if hot extrusion is used to compact the powder, the extruded material may be slowly cooled so that the period of time spent between the gamma prime solvus temperature and about 2200F is sufficient to permit gamma prime phase growth at grain boundaries.
The material at this point in the process sequence has a uniform fine grain size and consists of grains of the gamma phase which contain very fine gamma prime particles on the order of .5-3 microns and larger gamma prime phase particles on the order of 5-20 microns are located at the grain boundaries. This microstructure is shown in Fig. 2. The light colored phase is the gamma prime phase. The fine grain size of ~his material is stable at temperatures up to about the gamma prime solvus temperature. More importantly, this fine grain material displays superplastic behavior and may be ~3~i905 deformed at temperatures from about 1800F to the gamma prime solvus temperature with significant ductility and a low flow stress.
To illustrate this behavior, a one-inch bar of extruded RSR 104 material (composition shown in Table I) was heat treated for 4 hours at 2300F and then hot rolled. Several samples were hot rolled at temperatures oE 2200 and 2300F. The samples were red~ced from a starting thickness of 1.0 inch to a final thickness of .030 inch. Between 12 to 15 passes were employed to achieve the total reduction with a reduction per pass of from .030 to .080 inch. After each pass, the material was returned to a furnace maintained at 2200-2300F for a period of time sufficient to bring the material tempera-ture back up to the starting temperature. No difficulty was encountered in hot rolling this material and this success is somewhat surprising in view of the notorious difficulties encountered in hot working conventional superalloys. This behavior has important commercial ~0 implications since it will make possible fabrication of intricate shapes at low cost. The deformation step is of course optional and not required.
The next step is a directional grain growth step which makes use of the phenomena known as secondary recrystallization. An apparatus is required which can produce a large temperature difference over a small distance (a steep gradient). The hot end of the gradient `` ` 1~3~905 exceeds t~e ga~a pri~.e solvus t~,perature DUt iS less th2n the incipient melting te~perature. Th~ article is sLo~71y ~oved relative to the thermal gr2dient so that it is progressively and selectively heated to ~ te~.perature bet~,een the ga~.a primQ sol~Jus temperatùre and the incipient r,elting temperature. As the article is heated within this temperature range, the ga~.a pri~e particl~s which have stabilized the grain boundarie~ dissolve permitting grain gro~th to occur. Grain grot~th occurs, the driving force being a reduction in grain bou~dary area. The grains will grow in the direction of relat.ive motion between the article and the ther~al gradient and the resultant grains will be elongated having a major dimJension ~7hich is at least ten times the minor dimension.
The thermal gradient should be as steep as possiDle in order to ~aximiz2 the directiona,lit~- o~ the grains and mini~i~e the possibi'ity of laterzl grain gro~.7th and resultant transverse grain boundaries. The ther~al gradient should be at least 20F per inc'n ~.ea~ured at the ga~lm2 prime solvus, but prererably gr~ater. The experimenta~ work reported herein was done with a thermal 0~
/; ,,,.-j,fgradient of 150 to 300F per inch1. The Oradient T~as obtained by using induction heatinD with a graphit2 susceptor through ~hich the article passes The raLe at which the article c2n be ~oved rel2tive -o the ther~al gradient 2ppears to be on the order o~ 2bcut one-hal inch per hour. For obvious co~.~.erci~l re~sons, it is 3 ~
~36905 desirable to move the workpiece relative to the gradient at as rapid a rate as possible.
Following the directional secondary recrystallization step, the material will have a rather coarse gamma prime microstructure, within the elongated grains, as shown in Fig. 3. This is a consequence of the slow cooling rate through the gamma prime solvus temperature which permits a large amount of gamma prime phase nucleation and growth. It is preferred that as a final step the entire article be solution heat treated at a temperature between the gamma prime solvus temperature and the incipient melting temperature for a period of time suffi-cient to produce a solid solution, then rapidly cooled and aged to produce a refined gamma prime structure.
The conditions experimentally used were 4 hours above the gamma prime solvus followed by an air cool to room temperature and the article was then reheated to 1950F
for 4 hours and 1600F for 12 hours. Those skilled in the art will appreciate that these conditions may be ~0 widely varied and that they may be combined with other operations such as coating operations which involve heating the article. The effect of the heat treatment is to produce a very refined dispersion of gamma prime part;cles. Fig. 4 shows the microstructure after the aging step. The refined microstructure contributes to the improved mechanical properties.
Table I gives the compositions of three experimental 1~3~905 alloys, processed according to the invention, denoted as RSR alloys and two commercial prior art superalloys, PWA 1422 and PWA 1419. RSR 104 alloy is the basic ternary Ni-Al-Mo. RSR 166 is the same alloy composition as RSR 104 with the addition of 150 ppm Y. RSR 143 is a Ni-Al-Mo alloy with 6% Ta added. PWA 1422 is one of the strongest commercially used nickel base superalloys.
PWA 1419 is a prior art nickel base superalloy containing Ni, Mo, Al and C. This alloy is claimed in U. S. Patent 3,655,462. Both PWA 1422 and 1419 are used in the directional solidified form as described in U. S. Patent 3,260,505. Directional solidification is a casting pro-cess which produces a microstructure containing elongated grains. The properties discussed below were obtained from directionally solidified samples. Fig. 5 is a graph which shows the creep properties of the various superalloys in Table I. Fig. 5 shows the time required to reach 1% elongation in creep as a function of stress and temperature. Time and temperature are combined in
2~ the form of the Larson-Miller parameter, where T is absolute temperature and t is time in hours. For example, if one wished to determine the stress level required to produce 1% creep deformation at 1800F in 100 hours, one would determine the Larson-Miller parameter to be 49.7.
Then referring to Fig. 5 it can be seen that the stress level required for RSR 143 material exceeds 40 ksi while the stress level required for P~A 1422 material, a very ~1136905 ~,1 rd æ ~
~ I I ~
r~ I I I C~l I
I I I ~ I
E-l I ~D I I I
H C) I I 1 ~1 ~ O O
¢ ~ ,,, ,_1, ~ `
Then referring to Fig. 5 it can be seen that the stress level required for RSR 143 material exceeds 40 ksi while the stress level required for P~A 1422 material, a very ~1136905 ~,1 rd æ ~
~ I I ~
r~ I I I C~l I
I I I ~ I
E-l I ~D I I I
H C) I I 1 ~1 ~ O O
¢ ~ ,,, ,_1, ~ `
3 I I 1 ~1 ~') .,1 E~ I I I C`J
oO
¢ oo Ul o~
h C~ I I I ~
O O
~7 O
1136~05 strong prior art superalloy, is only about 25 ksi. Even more interesting is that the stress required to produce 1% creep in PWA 1419 is less than 20 ksi. Since PWA
1419 has a composition whicX is very similar to the RSR
alloys, this figure demonstrates the significance of the novel processing steps employed. The difference between the curves for RSR 104 (no Ta) and RSR 143 (6% Ta) demonstrates the effect of Ta additions and the effect is seen to be most pronounced at high values of the Larson-Miller para~leter, corresponding to long exposure times and/or high exposure temperatures.
Fig. 6 shows creep deformation behavior of the RSR
alloys and PW~ 1422 presented in a somewhat different fashion. The curves show elongation as a function of time at 1900F with an applied load of 14.2 ksi. (The curve for RSR 143 is extrapolated from data obtained at 1900F and 30 ksi.) It can be seen that P~dA 1422 reaches 1% creep in about 100 hours and fails in about 250 hours.
The basic ternary RSR alloy, RSR 104, requires about 700 hours to reach 1% creep and fails in about 900 hours, a signiicant improvement over PWA 1422. A dramatic increase in creep li~e is seen to result when 6% Ta is added to the basic alloy and this is shown as RSR 143 which requires more than 16,000 hours to reach 1% creep and somewhat more than 17,000 hours to fail. These creep results are truly outstanding when compared with the prior art superalloys, however, both RSR 104 and 1~L3690S
RSR 143 show somewhat low total creep ductilities. The addition of 150 ppm of Y to the basic RSR 104 composition results in alloy RSR 166 and the addition of this small amount of Y raises the total creep ductility from about 3.3% to about 9.3%, a very significant increase. It is believed that this improvement by the addition of Y
would be generally observed throughout the composition ranges previously described.
Further evidence for the beneficial effect oE Ta on the microstructural stability of the present alloys can be seen in Figs. 7, 8 and 9. Fig. 7 shows alloy RSR 144 (composition similar to alloy RSR 143 but containing 3%
Ta and 7% Al) after a solution treatment at 2400F and a rapid cooling step. The grain boundaries are seen to contain small amounts of the light colored gamma prime phase. The microstructure of RSR 143 after a similar solution treatment is very similar. Fig. 8 shows the microstructure of RSR 144 solution treated at 2400F
and aged for 50 hours at 2000F. Massive amounts of the gamma prime phase can be seen at the grain boundaries after this aging step and it is believed that this pre-cipitation of the gamma prime phase at the grain bound-aries adversely affects the high temperature mechanical properties. Fig. 9 shows the microstructure of RSR 143 after a 2400F solution treatment and 50 hours at 2000F.
The difference between Fig. 8 and Fig. 9 is dramatic since the massive gamma prime phase which is evident in Fig. 8 is completely absent from Fig. 9. Thus, it is evident that a Ta level somewhere between 3 and 6 weight percent, e.g. 4%, is effective in suppressing precipita-tion of the gamma prime phase at the grain boundaries in the alloys of the invention.
Although this invention has been shown and described with respect to a preferred embodiment thereof, it should be understood by those skilled in the art that various changes and omissions in the form and detail thereof may be made therein ~ithout departing from the spirit and scope of the invention.
oO
¢ oo Ul o~
h C~ I I I ~
O O
~7 O
1136~05 strong prior art superalloy, is only about 25 ksi. Even more interesting is that the stress required to produce 1% creep in PWA 1419 is less than 20 ksi. Since PWA
1419 has a composition whicX is very similar to the RSR
alloys, this figure demonstrates the significance of the novel processing steps employed. The difference between the curves for RSR 104 (no Ta) and RSR 143 (6% Ta) demonstrates the effect of Ta additions and the effect is seen to be most pronounced at high values of the Larson-Miller para~leter, corresponding to long exposure times and/or high exposure temperatures.
Fig. 6 shows creep deformation behavior of the RSR
alloys and PW~ 1422 presented in a somewhat different fashion. The curves show elongation as a function of time at 1900F with an applied load of 14.2 ksi. (The curve for RSR 143 is extrapolated from data obtained at 1900F and 30 ksi.) It can be seen that P~dA 1422 reaches 1% creep in about 100 hours and fails in about 250 hours.
The basic ternary RSR alloy, RSR 104, requires about 700 hours to reach 1% creep and fails in about 900 hours, a signiicant improvement over PWA 1422. A dramatic increase in creep li~e is seen to result when 6% Ta is added to the basic alloy and this is shown as RSR 143 which requires more than 16,000 hours to reach 1% creep and somewhat more than 17,000 hours to fail. These creep results are truly outstanding when compared with the prior art superalloys, however, both RSR 104 and 1~L3690S
RSR 143 show somewhat low total creep ductilities. The addition of 150 ppm of Y to the basic RSR 104 composition results in alloy RSR 166 and the addition of this small amount of Y raises the total creep ductility from about 3.3% to about 9.3%, a very significant increase. It is believed that this improvement by the addition of Y
would be generally observed throughout the composition ranges previously described.
Further evidence for the beneficial effect oE Ta on the microstructural stability of the present alloys can be seen in Figs. 7, 8 and 9. Fig. 7 shows alloy RSR 144 (composition similar to alloy RSR 143 but containing 3%
Ta and 7% Al) after a solution treatment at 2400F and a rapid cooling step. The grain boundaries are seen to contain small amounts of the light colored gamma prime phase. The microstructure of RSR 143 after a similar solution treatment is very similar. Fig. 8 shows the microstructure of RSR 144 solution treated at 2400F
and aged for 50 hours at 2000F. Massive amounts of the gamma prime phase can be seen at the grain boundaries after this aging step and it is believed that this pre-cipitation of the gamma prime phase at the grain bound-aries adversely affects the high temperature mechanical properties. Fig. 9 shows the microstructure of RSR 143 after a 2400F solution treatment and 50 hours at 2000F.
The difference between Fig. 8 and Fig. 9 is dramatic since the massive gamma prime phase which is evident in Fig. 8 is completely absent from Fig. 9. Thus, it is evident that a Ta level somewhere between 3 and 6 weight percent, e.g. 4%, is effective in suppressing precipita-tion of the gamma prime phase at the grain boundaries in the alloys of the invention.
Although this invention has been shown and described with respect to a preferred embodiment thereof, it should be understood by those skilled in the art that various changes and omissions in the form and detail thereof may be made therein ~ithout departing from the spirit and scope of the invention.
Claims (11)
1. A method for producing a high strength aligned grain superalloy article from an alloy which contains from 5-10% by weight Al, from 8-21% by weight Mo, zero to 12% by weight Ta, zero to 1% by weight Y, zero to .1% by weight C, zero to .05% by weight B, zero to .1% by weight Zr, zero to 12% by weight Cb, zero to 12% by weight Ti, zero to 12% by weight W, zero to 5% by weight Cr, zero to 5% by weight Co, balance nickel, including the steps of:
a. forming the alloy into a powder;
b. compacting the powder alloy into an article;
c. heating the article at a temperature between 2200°F
for a period of time in excess of one hour, and the gamma prime solvus so as to form gamma prime particles at the grain boundaries;
d. progressively and selectively heating the article by causing relative motion between the article and a thermal gradient, said thermal gradient having a hot end temperature which lies between the gamma prime solvus temperature and the incipient melting temperature, so that the gamma prime particles at the grain boundaries dissolve and directional grain growth occurs in the direction of relative motion between the article and the thermal gradient;
whereby said superalloy article has an aligned grain structure and at a temperature of 1800°F a stress in excess of 33 ksi is required to produce more than 1% of elongation in a time period of 100 hours.
a. forming the alloy into a powder;
b. compacting the powder alloy into an article;
c. heating the article at a temperature between 2200°F
for a period of time in excess of one hour, and the gamma prime solvus so as to form gamma prime particles at the grain boundaries;
d. progressively and selectively heating the article by causing relative motion between the article and a thermal gradient, said thermal gradient having a hot end temperature which lies between the gamma prime solvus temperature and the incipient melting temperature, so that the gamma prime particles at the grain boundaries dissolve and directional grain growth occurs in the direction of relative motion between the article and the thermal gradient;
whereby said superalloy article has an aligned grain structure and at a temperature of 1800°F a stress in excess of 33 ksi is required to produce more than 1% of elongation in a time period of 100 hours.
2. A method as in claim 1 wherein the article is deformed at a temperature between about 1800°F and the gamma prime solvus temperature between steps c. and d.
3. A method as in claim 1 wherein subsequent to step d. the article is solution treated at a temperature between the gamma prime solvus and the incipient melting temperature, then rapidly cooled and aged to provide a refined gamma prime microstructure.
4. A method as in claim 1 wherein the powder is formed from molten metal by a high rate solidification process.
5. A method as in claim 1 wherein the powder is compacted by hot extrusion.
6. A method as in claim 1 wherein the alloy contains up to about 12% by weight Ta as a partial equiatomic replacement for Al.
7. A method as in claim 1 wherein the alloy contains up to 1% by weight Y, up to .1% by weight C, up to .05% by weight B, up to .1% by weight Zr, and mixtures thereof.
8. A method as in claim 6 wherein part of the Ta is replaced by an equiatomic amount of Cb, Ti, or W, or mixtures thereof.
9. A method as in claim 1 wherein part of the Mo is replaced by W.
10. A method as in claim 1 wherein up to 5% by weight Cr, up to 5% by weight Co or mixtures thereof are added as partial replacement for the Ni.
11. A wrought recrystallized powder metallurgy article which comprises elongated aligned grains of the gamma phase which contain a fine dispersion of the gamma prime phase, said article having a composition of 5-10% by weight Al, 8-21% by weight Mo, up to 12% by weight Ta, up to 1% by weight Y, balance essentially Ni, said article being capable of withstanding stresses of up to 33 ksi at 1800°F for 100 hours without undergoing more than 1% elongation.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| US91315578A | 1978-06-06 | 1978-06-06 | |
| US913,155 | 1978-06-06 |
Publications (1)
| Publication Number | Publication Date |
|---|---|
| CA1136905A true CA1136905A (en) | 1982-12-07 |
Family
ID=25432976
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA000324088A Expired CA1136905A (en) | 1978-06-06 | 1979-03-22 | Superalloy composition and process |
Country Status (2)
| Country | Link |
|---|---|
| AU (1) | AU531066B2 (en) |
| CA (1) | CA1136905A (en) |
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| EP3290135A1 (en) * | 2016-08-29 | 2018-03-07 | Honeywell International Inc. | Methods for directionally recrystallizing additively-manufactured metallic articles by heat treatment with a gradient furnace |
Families Citing this family (2)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| GR80048B (en) * | 1983-12-27 | 1984-11-30 | Gen Electric | Yttrium and yttrium-silicon bearing nickel-based superalloys especially useful as comptible coatings for advanced superalloys |
| US4669212A (en) * | 1984-10-29 | 1987-06-02 | General Electric Company | Gun barrel for use at high temperature |
-
1979
- 1979-03-22 CA CA000324088A patent/CA1136905A/en not_active Expired
- 1979-04-03 AU AU45759/79A patent/AU531066B2/en not_active Ceased
Cited By (2)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| EP3290135A1 (en) * | 2016-08-29 | 2018-03-07 | Honeywell International Inc. | Methods for directionally recrystallizing additively-manufactured metallic articles by heat treatment with a gradient furnace |
| US10252337B2 (en) | 2016-08-29 | 2019-04-09 | Honeywell International Inc. | Methods for directionally recrystallizing additively-manufactured metallic articles by heat treatment with a gradient furnace |
Also Published As
| Publication number | Publication date |
|---|---|
| AU4575979A (en) | 1979-12-13 |
| AU531066B2 (en) | 1983-08-11 |
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